*3.2. Influence of DLC Coatings on the Transformation of the Sintered Ceramic Blanks Characteristics*

Figure 12 shows SEM images of the microstructure of thin sections of an experimental sample of 80% (90α10β) + 20% TiN ceramics (Figure 12a) and specimen of commercial ceramics (Figure 12b) after DLC coating deposition under identical technological conditions. Figure 12c presents an SEM image of a DLC coating's surface structure, and Figure 12d shows a TEM image of its volume structure. The coating thickness in total was 3.6 μm, including the thickness of 1.7 μm of the (CrAlSi)N sublayer and 1.9 μm of the functional DLC layer. It can be seen that the nitride sublayer has a columnar structure, while the outer DLC layer is characterized by an amorphous structure that microscopic methods cannot differentiate. The DLC layer does not have visible boundaries between grains in thickness, which are traditionally formed during the deposition of nitride coatings, while the surface of the DLC layer is a specific relief in the form of intergrown spherical segments of 0.2–1.5 μm in size. It should be noted that the formation of a characteristic relief is not affected by the ceramic base's surface layer's state before the coating deposition. The selected area diffraction (SAED) of the complex coating (Figure 12e,f) presents a complex phase of (CrAlSi)N sublayer and amorphous structure of DLC layers.

**Figure 12.** Microstructure of ceramic specimens with DLC coatings: (**a**) SEM image of a thin section of an experimental specimen 80% (90α10β) + 20% TiN ceramics with DLC coating; (**b**) SEM image of a thin section of commercial ceramics specimen with DLC coating; (**c**) SEM image of the DLC coating surface structure; (**d**) TEM image of the a two-layer DLC coating structure; (**e**) SAED of the (CrAlSi)N sublayer; and (**f**) SAED of the functional DLC layer.

Figure 13 shows the surface topographies of 80% (90α10β) + 20% TiN ceramic specimens before (Figure 13a) and after DLC coating deposition (Figure 13b). It is seen (Figure 13a) that the surface of the ceramic specimen includes a pronounced network of abrasive scratches, which are traditional for ceramic tools that are subjected to diamond sharpening. Besides, many craters from grains torn out during grinding are found on the ceramic surface previously shown in a series of previous works [17,21,25]. The noted features are defects of abrasive processing in essence, which are inherent in the process physics. The danger lies in the fact that ceramics are very sensitive to structural defects, which serve as stress concentrators and are often sources of tool's working surfaces spalling and chipping. It is possible to minimize or eliminate such defects in additional operations of finishing and polishing, which significantly increases the labor intensity and manufacturing cost of a ceramic product. The topography of the DLC-coated ceramic specimen surface (Figure 13b) demonstrates that the coating significantly modifies a thin surface layer's relief and significantly affects the size and shape of surface microroughness formed during the diamond sharpening stage. The coating fills the micro-grooves on the surface, thereby providing a kind of "healing" of the surface (it should be noted that a similar effect is observed when applying coatings of various compositions). There is reason to expect

that the coating will affect the characteristics of ceramic samples taking into account the described effect and the fact that the hardness of the formed DLC coating is not less than 25 GPa [73–77], which significantly exceeds the initial hardness of the tool ceramics, and the coating has a lower coefficient of friction and good strength of the adhesive bond with ceramic basis [78–80].

**Figure 13.** Topography of 80% (90α10β) + 20% TiN ceramic specimens surface areas: (**a**) before DLC coating deposition and (**b**) after DLC coating deposition.

Figure 14 shows histograms that illustrate the relationship between the applied critical load, at which destruction occurs, the path of punch movement along the z-axis until the destruction of the sintered ceramic samples, depending on the version of the powder composition. Data are given for samples before and after DLC coating deposition (quantification for each option was carried out based on four tests' results). The experimental data presented show that lower values of forces are required for the fracture of ceramic specimens 80% (70α30β) + 20% TiN and 90% (90α10β) + 10% TiN. It indirectly indicates their lower resistance to brittle fracture (this is confirmed by the studies presented in Figure 9). The coating has minimal effect on the increase in mean breaking force for these specimens. A different picture is observed for a sample of 80% (90α10β) + 20% TiN. Firstly, the average critical force is 5.2 kN, which is 10–25% higher than the corresponding indicator for the other two compositions. Secondly, the DLC coating deposition up to 6.8 kN (by 30%) increases the average critical force at which the sintered ceramic specimen breaks down (an increase in the punch travel path after which fracture occurs is also observed). The results obtained allow concluding that the coating can somewhat improve the fracture resistance of sintered ceramics based on SiAlON due to the transformation of the surface properties. However, it is only observed if the coating is applied to a ceramic base with a satisfactory combination of crack resistance and strength. It can be assumed that the improvements are a consequence of the effect mentioned above of "healing" of surface defects to a certain extent.

A series of tribological tests were carried out at a load of 1 N and a sliding speed of 10 cm/s according to the "ball-ceramic disk" scheme (due to the nickel alloy's increased wear counter body, an Al2O3 ball was used as a material in these tests) to assess the contribution of the coating to the change in the properties of the ceramics based on SiAlON (80% (90α10β) + 20% TiN) surface layer. Figure 15 systematized the tests' results at a friction length of 200 m before and after DLC coating deposition during tests without thermal exposure and under high-temperature heating conditions. The traditional character of the change in the friction coefficients of ceramic samples over time is observed in the room temperature results. At the very beginning of the tests, the uncoated ceramic sample has a friction coefficient of 0.2, which increases rather quickly and reaches 0.8 after 100 m of distance, which remains until the end of the tests. A very stable behavior of the coating during the entire test cycle, when the friction coefficient was invariably at the level of 0.1, is observed after DLC coating deposition. This coating belongs to the anti-friction class and is traditionally characterized by a reduced coefficient of friction. Considering that a ceramic cutter during the processing of a nickel alloy is subjected to high thermal effects, which has little in common with operating products at room temperature, the results of tribological tests when heated to high temperatures are of the most significant interest. Under these conditions, SiAlON (80% (90α10β) + 20% TiN) samples without coating show unstable results, when the friction coefficient changes abruptly. Friction coefficient μ intensively increases after a value of 0.2 at the beginning and reaches a value of more than 0.7 after 25 m, then decreases to 0.45, and then it increases, decreases, and again increases, reaching more than 0.9 by the end of the tests. The authors described similar unstable results of other tool ceramics' behaviors (based on Al2O3) in a previously published study [25]. The DLC coating deposition onto ceramic samples strongly changes the conditions of frictional interaction: μ remains at a low level and varies slightly within the range of 0.09–0.15 for a sufficiently long time, and it begins to increase only after passing a distance of 150 m, reaching a value of 0.72 by the end of the tests. It can be assumed that the aforementioned is a consequence of the unique properties that a two-layer coating (CrAlSi)N/DLC possesses upon contact with the counter body under conditions of intense heat exposure. Superficial carbon DLC coating layers show poor results at increased thermal loads and often lose their initial microhardness, known from several authoritative works [12].

**Figure 14.** The relationship between the critical (breaking) load, the path of punch movement until the fracture of ceramic samples sintered from various powder compositions, where (1) is 80% (90α10β) + 20% TiN, (2) is for 90% (90α10β) + 10% TiN, and (3) is for 80% (70α30β) + 20% TiN before and after DLC coating deposition.

**Figure 15.** Dependences of the friction coefficient on the surface of experimental specimens made of SiAlON (80% (90α10β) + 20% TiN) ceramics on the friction path before and after DLC coating deposition under various test temperature conditions: without heating (20 ◦C) and with heating to 800 ◦C.

Nevertheless, due to the introduction of silicon into its composition during the deposition process, a formed DLC-Si layer makes it possible to significantly increase the DLC layer's thermal stability and expand the field of application of such coatings in this case [75,81,82]. In addition, the presence of a nitride sublayer under the DLC layer based on the Cr-Al-Si system increases the strength of the external DLC layer's adhesive bond with the tool base [83–85] and contributes to the formation of secondary wear-resistant phases during high-temperature heating [86–90]. When heated in oxygen, the coating components can form nonstoichiometric oxide phases, contributing to the change in the contact interaction between the ceramic article and the counter body nature [91–93].
